Rolled metal articles of high biaxial ultimate strength and production thereof



Aug. 16, 1966 MB, VQRDAHL 3,266,951

ROLLED METAL ARTICLES OF HIGH BIAXIAL ULTIMATE STRENGTH AND PRODUCTION THEREOF Filed March 12, 1965 5 Sheets-Sheet 1 INVENTOR ATTO R N EYS Aug. 16, 1966 M. B. VORDAHL 3,266,951

ROLLED METAL ARTICLES OF HIGH BIAXIAL ULTIMATE GTH A 0 ND PRODUCTI N THEREOF 5 Sheets Filed March'l2, 1965 -Sheet 2 INVENTOR ML TON B. V/PaqHL.

MWZ YhM ATTORNEYS.

L 3,266,951 BIAXIAL ULTIMATE UCIION THEREOF 1966 M. B. VORDAH ROLLED METAL ARTICLES OF HIGH STRENGTH AND PROD Filed March 12 1965 5 Sheets-Sheet 3 INVENTOR v MLTONB VcBPDAHL.

J ATTORNEY5.

United States Patent corporation of New Jersey Filed Mar. 12, 1965, Ser. No. 445,824 5 Claims. (Cl. 148-32) This application is a continuation-in-part of my co pending application Serial No. 31,949, filed May 26, 1960, now abandoned.

This invention pertains to rolled metal products and the production thereof characterized in having a bilateral utilmate strength greatly in excess of the ultimate strength as stressed in a single direction.

The conventional method of determining the ultimate strength of a metal is to form a tensile test bar or flat having an intermediate length of restricted cross-section as compared to the ends, and to grip the ends in a tensile test machine adapted to pull the ends in opposite directions until the specimen ruptures. For a normal test the specimen ruptures in the intermediate portion of restricted section; otherwise it is discarded and another specimen tested. The ultimate strength is the ratio of the load producing the rupture divided by the original cross-sectional area at the point of rupture. Thus the conventional method of measuring the ultimate strength of the material is to tension stress it in one direction only to the point of rupture.

Another way in which the ultimate strength of the metal could be measured would consist in clam-ping a sheet of the material in a circular ring clam-p, like the skin of a drum, and bulging it with hydraulic pressure exerted on one side until the sheet ruptured. A crude analogy to this testing technique would be to introduce air under pressure into the interior of a kettle-drum and gradually increase the pressure until the skin burst. As the pressure is increased, the sheet material bulges out in accordance with a substantially spherical surface, the curvature of which increases with the pressure up to the bursting point. Since the fluid, pressure is uniform over the spherical bulge area of the sheet, any small circular area thereof will be more or less uniformly tensioned in all directions about its center, referred to herein as bilateral stressing. The b-ilateral ultimate strength at the bursting load can be computed in conventional manner, as discussed below, from the bursting pressure, the orig inal thickness of the sheet 'and the radius of spherical curvature of the material at the bursting pressure.

My investigations have shown that when pure metals or alloys are rolled into sheet by conventional methods and tested in this way for bilateral ultimate strength and also tested in the conventional way for ultimate strength in the manner first above described, that in general there is a close correspondence between the two ultimate strength determinations thus made. That is to say, the bilateral ultimate strength will in general not exceed the ultimate strength as conventionally determined by more than about -20% at the utmost.

Now I have discovered in accordance with the basic concept of the present invention, that if titanium and alloys thereof which have at normal or atmospheric temperatures a closeapacked hexagonal crystal structure are processed by heating and extensive rolling into sheet, i.e., to the extent of at least 75% total reduction as a result of the rolling operation, and in the temperature field in which the hexagonal crystal structure is retained or predominates, the bilateral ultimate strength of the material so processed greatly exceeds its uni-directional ultimate strength, and to the extent of at least 25%.

The materials in which this effect is found to be most pronounced are the titanium-base alloys which contain substantial amounts of aluminum, for example about 28% aluminum and wherein the balance of the alloys is such as to imp-art at normal or atmospheric temperatures, a crystal structure which is wholly or at least predominantly of the close-packed hexagonal type, referred to herein as all-alpha or a predominantly alpha crystal structure, respectively. By a fipredominantly alpha crystal structure, I means a crystal structure wherein the alpha phase accounts for at least 90% by volume of the metal. As above stated, the requirements for producing the enhanced bilateral strength as compared to the uni-directional strength is that the metal or alloy shall be rolled to the extent of producing at least total reduction, and that the rolling shall be accomplished at a temperature in which the close-packed hexagonal or alpha crystal structure predominates, such reduction not to be followed by any working at a higher temperature or by any heating to a higher temperature.

As illustrative examples, I have found that if commercial purity titanium having an ultimate strength of about 90,000 p.s.i. as conventionally measured, is rolled into sheet to the extent of at least 75% total reduction, and then annealed, both the rolling and annealing being at temperature below the beta transus, i.e., in the all-alpha temperature field, and burst tested in the manner aforesaid, its bilateral ultimate strength as thus determined, is about 120,000 p.s.i., i.e., about 35% greater than its unidirectional ultimate strength. A more striking example is provided by the all-alpha titaniumbase alloy Ti-7Al-12Zr, having a unidirectional ultimate strength as annealed of about 135,000 p.s.i. but a bilateral ultimate strength when rolled and annealed as above, of about 250,000 p.s.i., an increase of about As discussed below, both titanium metal and the aforesaid titaniumaluminum-zirconium a1- loy have a close-packed, hexagonal crystal structure at normal or atmospheric temperatures, and also at temperatures below the beta transus of about 885 C. or 1625 F. for the pure titanium metal and higher than that for the aforesaid alloy.

Without binding myself to any particular theory in explanation of the aforesaid extremely surprising results, my investigations indicate that the explanation resides in the manner in which the close-packed, hexagonal crystals are oriented in the sheet material produced as above. It appears that extensive rolling of the material in the temperature range in which the hexagonal crystal structure is retained, imparts thereto a preferred orientation such that the basal planes of these crystals lie in the surface planes of the sheet or parallel thereto.

Development of a preferred orientation with basal planes parallel to the rolling plane depends, of course, on obtaining extensive basal slip. The most active slip planes tend to rotate into parallelism with the rolling plane, with slip direction parallel to rolling direction. With a metal like titanium or zirconium or ran all-alpha alloy thereof, prismatic slip apparently occurs first, with pyramidal and basal slip occurring later, and with extensive basal slip occurring after extension rolling, and with the basal texture thus appearing importantly only after extensive rolling such as to provide about 75 total reduction. Pyramidal slip and, most importantly, twinning (which occurs on pyramidal planes only) tend to cause some basal scatter. Aluminum inhibits twinning in titanium, and thus tends to promote development of the greatest degree of parallelism between basal planes and rolling planes, and in turn to produce the greatest effect in increasing biaxial strength.

If, therefore, an elongated, rectangular strip specimen having its crystals oriented as aforesaid, is stressed in tension in the direction of its length, the specimen will deform by lengthening and narrowing, by slippage of the crystals relative to each other along their paralleldisposed prism faces or rectangular atom planes extending from top to base, parallel to the short axis of each. This slippage results from the fact that the tensioning stress is applied to the crystal array only in one direction, namely, that extending lengthwise of the specimen.

In contrast, however, if the specimen, is bilaterally stressed, the crystals are restrained from slipping relative to each other in the above described manner, since uniform tensioning is now applied to each crystal in all radial directions thereabout, so that it is equally restrained against slipping in any radial direction. And since under the aforesaid preferred-orientation assumption, the crystals in any one layer of the lattice are attached lengthwise on those of the layer below, the bilateral stressing also prevents any thinning deformation of the sheet material in response thereto, since there is no tendency of the stacked layers to slide on their basal planes with respect to each other. The only other planes along which the crystals might tend to slip relative to each other, are higher index prism planes, or the pyramidal planes extending obliquely to the crystal axes between atom pairs at the top and base. But slip on any of the three low index planes tends strongly to be only in directions parallel to the short sides of the prism faces (directions S and S in FIGURE 1). Thus it is apparent that the thinning necessary to permit are-a extension during bilateral deformation becomes increasingly difficult to accommodate by slip as the basal orientation becomes more nearly perfect or complete.

As a result, the material withstands bilateral stressing without appreciable deformation, more nearly to the point of crystal rupture, than it does in the necking-down deformation under uni-directional stress, and hence ruptures at a much lower applied load in the latter case than in the former. The fact that the experimental test results are in complete accord with those which theory predicts under the above preferred crystal orientation assumption, quite conclusively confirms this is What actually occurs as a result of rolling and annealing the material in the temperature range within which the close-packed, hexagonal crystal structure is retained.

This is also confirmed by other tests as follows. If two sheets of the aforesaid material are welded together, and bilaterally stressed to rupture, the rupture always occurs at the weld seam, the crystal structure of which approaches a random orientation. This effect does not depend on weld brittleness or weakness. My tests have shown that welds which failed under biaxial stress failed ductilely, and that their uniaxial strengths were nearly the same as that of the base metal.

Again, if a metal having at normal temperatures, a cubic crystal structure, such as the all-beta titaniumbase alloy Ti13V11Cr3Al, having a body-centered cubic crystal structure, is rolled into sheet, the ultimate strength as determined by biaxial stressing, is found to be substantially that obtained by the conventional unidirectional tensile test procedure. The explanation for this is that even assuming a preferred orientation of the cubic lattice structure such that surface planes of the cubic crystals lie in or parallel to the surface of the sheet, there are low index slip planes and directions that will accommodate thinning. Thus the sheet specimen can deform under bilateral stress in substantially the same manner as in uni-directional stress, by slippage along several inclined planes. Thus again theory predicts what experiment confirms, namely, that in metals having a cubic lattice, a preferred crystal orientation offers little advantage as regards bilateral versus unidirectional stressing.

Having thus described the invention in general terms,

reference will now be had for a more detailed description to the annexed drawings herein:

FIGURE 1 shows schematically, in perspective, the atomic spatial arrangement of a close-packed, hexagon-a1 crystal; and FIGURE 2 similarly shows that of a bodycentered cubic crystal.

FIGURE 3 is a view in plan, and FIGURE 4 is a view in perspective, on a greatly enlarged scale, of the lattice structure of a rolled sheet specimen of a metal having a hexagonal crystal structure wherein the crystal lattice has the preferred orientation above-mentioned in which the basal planes of the crytsals lie in or parallel to the surface of the sheet specimen. FIGURE 5 is a a plan view of the lattice structure of FIGURE 3 as distorted by uni-directional tensioning stress. Actually, of course, the deformation on a macro scale involves much grain fragmentation, distortion and rotational accommodation of the movement. Nevertheless, the effect is about as shown.

FIGURE 6 is a view in plan of a rolled sheet specimen of a metal having a body-centered cubic latice structure, in which the cubic crystals are oriented with cubic faces lying in or parallel to the surface of the sheet. FIGURE 7 is a diagrammatic showing of the manner in which the sheet specimen of FIGURE 6' deforms under uni-directional stressing in tension. FIGURE 8 is a prospective view of a pair of stacked crystals of the FIGURE 6 lattice structure, and FIGURE 9 a view in perspective showing how the FIGURE 8 crystal lattice deforms under bilateral stressing in tension. These views are schematic only, and planes and directions of low index are shown for simplicity of illustration only. The fact that other planes and directions are actually involved does not impair the general validity of the illustrations.

FIGURE 10 is a view in axial section through an apparatus for biaxially stressing a sheet specimen to rupture in the manner above described.

Referring to the close-packed, hexagonal crystal structure of FIGURE 1, the slip of a metal having this crystal structure generally tends to occur most easily along the basal planes 1 and 2. In certain such metals, notably titanium, however, early slip occurs most readily along the prism planes, as at 3, and along the pyramidal planes, such as that shown by the shaded area 4, and perhaps least along the basal planes 1, 2 until after a large amount of deformation. Then basal slip occurs increasingly, and, as with other hexagonal, close-packed metals, basal orientation very roughly parallel to the rolling plane begins to appear. With pure titanium, a moderate amount of angular scatter is evident. This scatter is greatly increased by the presence of beta stalbilizers, and decreased by aluminum. Working and/ or heating in the a'lpha beta tempenature range or above grossly increases the scatter, and may even substantially randomize the structure. Such randomization tends to make further rolling easier, of course, and is thus desirable according to standard practice.

Referring now to FIGURE 2, showing the body-centered cubic crystal structure, such as that of titanium above the beta transus and also at room temperature in titanium-base alloys containing a sufficient amount of beta-stabilizing elements, such as molybdenum, vanadium, columbium, tantalum, chromium, manganese, iron, etc., the slip planes are primarily the inclined planes, indicated by the shaded areas, as at 7 and 8, extending between oppositely disposed lat-om pairs in oppositely disposed cubic face surfaces, as shown. Several other planes in addition are usually active, however, during deformation of b, c, c metals.

Referring now to FIGURES 3 and 4, there is shown on a greatly magnified scale, a sheet specimen composed of a metal having a close-packed, hexagonal crystal structure, in which the basal planes of the crystals are oriented to the plane of the sheet and parallel thereto.

In plan view, FIGURE 3, the crystals will the arranged in contiguous fashion, as shown at F to Q, inc., and in depth, as shown in FIGURE 4, at a, b, c, d, etc.

Referring now to FIGURE 5, if the specimen of FIG- URE 3 is subjected to uni-directional tensile stress in the direction of the arrows 10, 11, the metal specimen will elongate in the direction of these arrows, and narrow in width, by slippage of the crystals along their prism planes, as at 12, 13 and 14, 15. Thus, the ultimate strength of the material as determined by unidirectional tensile stressing, as at 10, 11, will result from necking down of the specimen in the manner illustrated as the load is increased until the specimen ruptures, and the ultimate strength will be the applied road at rupture divided by the original cross-sectional area of the specimen.

In contrast, and referring to FIGURE 3, if the specimen therein shown is uniformly tensioned biaxially, as indicated by the arrows 16, 17 and 18, 19, there can be no slippage of the crystals F, G, H, etc., along their prism planes, as at 12, 13, and 14, 15, since each crystal, such as crystal 1, of the crystal lattice, is now subject to uniform tension stress in the directions 16, 17 and 18, 19, normal to each other.

Nor, referring to FIGURE 4, can the material deform under the bilateral stressing action by necking down in depth, since the crystals in each layer of the crystal lattice stand on end in each layer as at a, c, and are stacked on the similarly disposed next lower layer of crystals, such as b, d. Thus there is no tendency for the bilateral stressing in any radial direction, to cause the layers to slide on their basal planes with respect to each other, such as layer a, c, to slide with respect to layer b, d. Also, since the easy slip direction of the other planes lies in the basal plane, there is no low stress means, via slip, of accommodating thinning.

Referring now to FIGURE 6, which shows on a greatly magnified scale, a specimen of rolled sheet material composed of a body-centered, cubic metal, in which the cubic crystals are oriented with their faces in the plane of the sheet or parallel thereto, if the specimen is tensioned uni-directionally, as at 21, 22 of FIGURE 7, slip will occur along the diagonal slip planes, with results about as indicated in FIGURE 7. Thus, assuming the specimen to 'be subjected to uni-directional tension stress by applying tension in the direction of the arrows 21, 22, the specimen will elongate in the direction of the tensile stress and neck down in the direction normal thereto, by slippage of the crystal lattice along the diagonal slip planes as at 23, 24, in the manner illustrated in FIGURE 7.

If now, reverting to FIGURE 6, the specimen is subjected to biaxial, uniform tension stress, as at 21, 22 and 25, 26, there will be little tendency for slip to occur in the plane of the sheet, since each crystal, such as S, is subjected to uniform tension in all radial directions. in depth in the specimen in the manner illustrated in FIGURES 8 and 9. Thus, considering FIGURE 8 wherein two crystals A and B of the FIGURE 6 lattice are shown stacked one atop the other, slip can concurrently occur along the normally disposed slip plates 27 and 28 of crystals A and B when subjected to bilateral tension ing in the normally disposed directions 21, 22 and 25, 26. The manner in which this occurs is clearly shown in FIGURE 9, wherein the tensioning stress in the direction 25, 26 causes portion g of crystal B to slide along the diagonal slip plane 28 with respect to portion h of this crystal. Concurrently with this slippage, the tensioning stress applied in the direction 21, 22, i.e., normal to that of 25, 26, causes portion e of crystal A to slide along the diagonal slip plane 27 with respect to portion 7 of this crystal. It is again to be emphasized, of course, as previously noted, that these views are schematic only, and that planes and directions of low index are shown for simplicity of illustration only. The fact that other planes and directions may actually be involved does not, however, impair the general validity of the illustrated actions. Thus under the action of the bilateral loading, the specimen will become progressively thinner in depth and greater in area as the loading is increased to the rupture load of the specimen. The action in bilateral loading is, therefore, essentially the same as in uni-directional loading, since both produce a necking down or thinning of the specimen until rupture occurs. Theory thus predicts that the bilateral ultimate strength 'will not differ materially from the uni-directional ultimate strength for a metal having a cubic lattice structure. That tests confirm this is shown below.

Referring now to FIGURE 10, which shows in axial section a suitable apparatus for biaxially stressing to rupture or bulge testing rolled sheet metal blanks, the sheet specimen 40, which may be of hexagonal or circular shape, is clamped about its periphery, between a However, slip along diagonal planes can occur pair of dies 41, 42, of circular configuration in plan view. The lower die 42 is peripherally grooved adjacent to its inner edge as at 43, for seating a rib 44 extending peripherally about the inner edge of the upper die 41, with the sheet specimen 40 clamped therebetween, thus to crimp and anchor the sheet and lock it in position about its peripheral edge as shown. The upper and lower dies 41, 42, are seated in upper and lower housings 46, 47, respectively, which housings are bolted together, as by means of bolts 48, 49, having nuts threaded thereon such as 50, 51. In order to provide an oil seal between the die and sheet assembly, the rib 44 of the upper die 41 is grooved, for reception of an O-ring 45. This provides a sealed-in-cavity, as at 52, between the upper die 41 and the sheet specimen 40, into which oil under pressure is introduced through a pressure line 53. As the pressure is increased, the sheet specimen is thus bulged downwardly within a cavity 55, formed by central openings in the lower die 42 and lower housing 47.

The central opening of the lower die housing 47 is provided 'with an inner peripheral rib 57 on which is seated a flanged portion 58 of a substantially cylindrical housing 59, in which is rotatably mounted a spherometer 60 having a pair of contact elements 61 and 62, which bear against the underside of the sheet specimen 40 to provide a measurement of the bulge depth of a fixed chordal diameter. To this end, the outer contact ele ment 62 is mounted on an offset arm 63 integral with a sleeve member 64, which is rotatably mounted within the housing 59, so that the outer contact element 62 may be rotated about the axis of symmetry of the bulged specimen 40. The inner contact element 61 contacts the sheet specimen 40 at the center thereof, and is mounted on a rod 65 which is axially displacea'ble within the sleeve 64.

To the lower projecting end of the rod 64 is clamped, as at 66, a dial indicator 67, having an actuating arm 68, which is continually pressed by a spring of the dial indicator (not shown) against the lower side of a ring member 69, mounted on the lower projecting end of the sleeve 65. Thus relative depression of the central contact member 61 with respect to the chordal contact member 62, produces a reading on the dial indicator 67 corresponding to the bulge depth of the chordal segment. Thus the spherical radius of. curvature R of the bulge of the specimen may readily be calculated from the bulge depth h as read on meter 67, and the chordal diameter a ac cording to the formula R=a /4h. For safety reasons, the upper die 41 is fitted with a bleeder screw 70 which is used to bleed air out of the system prior to application of the hydraulic pressure. Also the housing 59 is provided with an axially extending bore 71 connected with an oil drain line 72 for draining off oil.

In conducting bulge tests on specimen 40 with the apparatus of FIGURE 11, four readings of bulge depth TABLE I.-SAMPLE DATA SHEET I 1% diameter. Burst Pressure 1,5

R5626. Solution treated and welded.

R=the mean radius of curvature of the bulge in inches. T=the original thickness of the sheet specimen in inches.

Using the above formula, the original thickness of the sheet specimen 40 is used to calculate the stress as in Dial Gage Calcu- Gage Readings Avg. R+'I/2, R, lated Pressure, Depth, Inches Inches Stress, (Depth) 2 p.s.i. Inches p.s.i.

Referring to the above table, the first column gives the test pressure of the oil. The next four columns under the heading Dial Gage Readings give the dial indicator readings for the 90 settings of contact 62 around the bulge. These values are averaged in the next column to give a mean bulge depth. This bulge-depth value is then used to obtain the average radius of curvature of the bulge surface. This radius is tabulated in the next column, headed R-l-T/Z. A mean radius of curvature R is entered in the next column. The biaxial stress of the sample may then be calculated using the following formula, which is recognized as the expression for stress in a thin-walled spherical shell subjected to internal pressure, it being readily shown that this formula is also valid for a spherical segment.

where S=the biaxial stress in p.s.i. (This stress is uniform on all great circles of the spherical bulge sample as sections remote from the clamped edge.)

P=pressure in p.s.i.

a conventional tensile test. The ultimate strengths obtained in this manner are more conservative than would be obtained if the thickness at each pressure were used. The calculated values of stress are tabulated in the ninth column of Table 1.

Using the apparatus of FIGURE 10 and the method of bulge testing the specimens above described with reference thereto, bulge tests to rupture were made on rolled sheet specimens which had undergone a rolling reduction in excess of 75% and which were made of titanium and various titanium-base alloys with results as set forth in the following Tables II-IV, inc. The tests recorded in Table II were made on sheet specimens of titanium and titanium-base aloys having an all-alpha or closepacked hexagonal crystal structure; those in Table III were made on titanium-base alloys having at normal temperatures an all-beta or body-centered cubic crystal structure; and those in Table IV were made on titaniumbase alloys having at normal temperatures a mixed alphabeta structure, consisting for the alpha phase of a closepacked hexagonal crystal structure and for the beta phase of a body-centered cubic crystal structure.

TABLE II.SUMMARY OF BULGE TESTS ON ALPHA TITANIUM ALLOYS Nomlnal Orig Ten- Tens. Biaxial Strength Average Composition, Condition Thick sile, Elong. Ult. Ratio, Biaxial Percent In. p.s.i.- Percent Stress, Col. 6 Elong Bal. T1 in 2 e p.s.i. b Col. 4 Percent 5Al-2.5Sn Annealed 056 132 16 1. 44 5Al-2.5Sn d 5Al-2.5Sn 5Al-2.5Sn 5Al-2.5Sn 5Al-2.5Sn

P.s.i. X 1,000 c Unalloyed titanium.

TABLE III.SUMMARY OF BULGE TESTS ON 13V-11Cr-3Al BETA TITANIUM ALLOY Biaxial Strength Average Orig Tensile, Tens. Ult. Ratio, Biaxial Condition Thick p.s.i., b percent Stress, Col. 5- Elong, In. in 2 a p.s.i. Col. 3 percent in 2 l 72 hrs. 900 F .041 200 8 Specimen cracked during tightening of die Sol. Trtd plus weldecL 040 150 1. 0 Sol. Trtd. plus welded 0425 150 1.0

Sol. Trtd. plus Welded 042 Failed at low pressure with no 60 hrs. 900 F. plus ductility. min. 1,100 F 042 S01. Trtd. plus welded" 042 151 2.0

Average of L and T directions. b P.s.i. X 1,000. o Elongation in 0.6 gage length.

TABLE IV.SUMMARY OF BULGE TESTS ON ALPHA-BETA TITANIUM ALLOYS Nominal Orig. Tensile Tens. Biaxial Strength Average Composition Thick, Elong. Ult. Ratio, Biaxial Percent Condition In. p.s.i. b Percent Stress, Col. 6+ Elong, Bal. Ti in 2" p.s.i. b 001. 4 Percent in 2 B Annealed 047 130 12 174 1. 34 3. 0 Annealed 050 130 12 170 1.30 3.0 1,700 F. min. WQ, plus 0495 161 8 d 130 81 Zero 2 hrs., 1,100 F. 1,700 F.30 min. WQ plus 0487 154 0 10 d 115 75 Zero 2 hrs., 1,200 F. 1,700 F.30 min. WQ plus 0305 178 c 5 d 142 80 Zero 2 hrs, 900 F. 1,700 F.30 min. WQ plus 0295 188 c 2. 5 d 143 76 Zero 2 hrs., 900 F. 6Al-4V Annealed and welded---" .048 163 1. 5 4Al3Mo-1V Sol. Trtd 043 146 13 162 1. 11 8.0 4Al-3Mo-1V Sol. Trtd 043 146 13 156 1.08 7. 0 4Al-3Mo-1V c Sol. Trtd. plus welded 043 130 1. 0 16V-2.5Al. Sol. Trtd 041 105 19 125 1.19 16V-2.5Al. Sol. Trtd 041 105 19 119 1. 14 16V-2.5Al Sol. Trtd. plus 960 F.-4 041 Sample cracked in tightening hrs. die. I I

11 Average of L and T directions. b P.s.i. X 1,000.

Elongation in 0.6 gage length. d Fractured at edges.

Referring to the test results in Table II for the tests conducted on rolled sheet specimens, having at room temperature the all-alpha or closepacked hexagonal crystal structure, it will be seen that all exhibited a considerably higher bursting or biaxial ultimate strength (Column 6), than would be expected from their ultimate strengths as normally determined in the conventional tensile test (Column 4). Thus the ultimate strengths as conventionally determined for the T i-5Al-2.5Sn alloy specimens ranged from about 132,000-135,000 p.s.i. (Column 4), as compared to biaxial strengths (Column 6) ranging from about 190,000-210,000 p.s.i., so that in Ibiaxial loading the material 'Was about -55% stronger than in unidirectional loading. For the Ti-5Al-12Zr alloy, the results are even more outstanding, this material being almost 100% stronger in biaxial loading than in conventional uni-directional loading as regards ultimate strength. It will further be noted that for the welded specimens, the ultimate strengths in lbiaxial loading were considerably less than for the non-welded specimens due, as explained above, to the more or less random crystal orientation in the welded portions. The biaxial strengths for the Welded specimens were, however, considerably higher than the ultimate strength of the alloy specimens as conventionally determined.

Cylindrical pressure vessels normally use dished ends, and the question therefore arose as to whether or not tests made on initially fiat blanks such as those of Table II, simulate this effect. To' investigate this condition, both types of tests were conducted on the Ti-5A1-2.5Sn alloy, in which the sheet specimens were pre-buldged and then annealed and re-test ed to failure by biaxial loading. The test results gave bursting test values of 190,000-200,000 p.s.i., which are in agreement with the data given in Table II for tests made on flat blanks. In these pre-bulge tests,

stress calculations were based on the thickness at the.

center of sample after the initial bulging.

Table III summarizes the bulge test results on Ti-13V- 11Cr-3Al alloy which has a beta or body-centered cubic structure at normal temperature. As noted in the table, the test results were run on sheet specimens in various heat-treated conditions, namely, as solution-treated and quenched and as solution-treated, quenched and thereafter aged for the times and at the temperatures indicated. Also welded specimens were included. It will be observed that although the solution-treated and quenched specimens exhibit excellent ductility, the ultimate strengths under biaxial loading approximates the normal ultimate strengths of the specimens, as conventionally determined. The double-aging treatment for 60 hours at 900 F. followed by hours at 1100 F. raised the strength and lowered the ductility, but with no appreciable advantage in 'biaxial as compared to uni-directional loading. The data. further show that the solution-treated welded specimens had about the same ultimate strength as the unwelded material as conventionally determined. Also that the aged materials were in general brittle under biaxial loading.

Referring to the test results of Table IV, for the alloy specimens having the mixed alpha-beta structure at normal temperatures, it will be seen that these show a moderate increase in strengthening under biaxial as compared to uni-directional loading in the annealed condition of the specimens, but were weaker in biaxial loading than in uni-directional loading in the solution-treated and aged conditions, due to the fact that edge fractures at relatively low stresses resulted wit-h no measurable ductility, although as normally tensile-tested, these samples showed tensile elongations of from 2:5 to 10%. In the Table IV alloys tested, the Ti-6Al-4V can be described as a near alpha" alloy, and the relatively high percentage of the alpha or close-packed hexagonal crystal structure phase probably account-s for the biaxial ultimate strength of 170,000 p.s.i. obtained in the annealed condition.

Comparison of the test data of Tables IIIV, conclusively establishes that whereas the ultimate strength of rolled sheet specimens of some metals and alloys having the hexagonal crystal structure is greatly increased under biaxial loading as compared to uni-directional loading, this is not true of the metals and alloys having the cubic crystal structure. Also in confirmation of the theories above discussed, it is further seen from a comparison of the data in Table IV for the mixed alpha-beta alloys with those of Tables II and III, respectively, for the allalpha and all-beta alloys, that the increase in biaxial strength appears to decrease progressively as the amount of the alpha or close-packed hexagonal crystal constituent is decreased in passing from the all-alpha to the all-beta or body-centered cubic structure.

The important titanium-base alloys having the closepacked, hexagonal crystal structure and within fa'bricable limits for the alloying additions of each, are those containing up to 50% zirconium, 23% tin, 18% antimony, cadmium, indium, zinc and silver, and 10% aluminum, the lower effective limit for each of these alloying additions being 0.05 to 1%, although normally the alloys will contain at least 2 to 3% in total amount of these alloying additions. Although aluminum has a facecentered cubic crystal structure and tin has a bodycentered tetragonal crystal structure, nevertheless when these elements are alloyed with titanium, the resulting crystal structure is that of alpha titanium, namely, closepaeked hexagonal.

The elements which impart the body-centered cubic crystal structure to titanium when alloyed therewith in sufilcient amounts, are vanadium, columbium, tantalum, molybdenum, chromium, manganese, iron, tungsten, copper, cobalt, nickel, silicon and beryllium, silicon and beryllium being however primarily compound formers and the same being generally true of copper, nickel and cobalt. As shown by the foregoing test results, increasing additions of these elements tend more and more to destroy the enhanced biaxial strengthening elfect of the resulting titanium-base alloy.

In order to maximize the prefer-red orientation of the alpha or close-packed hexagonal crystal lattice, the heat treating should be confined to the alpha temperature range and the same with respect to the rolling. Working by rolling at temperatures at which beta predominates is to be avoided. Where welding is required, directional chill- 12 ing of the weld tends favorably to orient the weld metal and to alleviate the loss in strength in biaxial loading, as does sharply limiting weld width.

As regards the titanium-base alloys having the allalpha or close-packed hexagonal crystal structure, the alpha-worked titanium-aluminum alloys show a higher degree of basal orientation as compared to other alpha alloys.

By the term titanium-base alloy as used in the appended claims is meant an alloy of titanium and one or more additional metals in which titanium constitutes at least 50% of the total weight.

I claim:

1. A pressure vessel composed at least in part of rolled metal sheet, the metal of which is selected from the group consisting of titanium and titanium-base alloys having at atmospheric temperatures predominantly a close-packed, hexagonal crystal structure, the crystals of which at least in the surface portions of said metal sheet have a preferred orientation with their basal planes in said surface portions of said sheet and parallel thereto, said metal sheet being characterized by a bi-axial ultimate strength at least 25% in excess of its unidirectional ultimate strength, said pressure vessel in its normal use subjecting said metal sheet to bi-axial tension stresses.

2. A pressure vessel composed at least in part of a rolled metal sheet, the metal of which comprises a titanium base alloy containing 1 to 50% of at least one alpha promoter selected from the group consisting of zirconium, tin, antimony, cadmium, indium, zinc, silver and aluminum, but not to exceed 23% tin, 18% antimony, 15% each of cadmium, indium, zinc and silver, and 10% aluminum, up to 10% of beta promoters, and the balance substantially titanium, characterized in having at atmospheric temperatures a predominantly alpha crystal structure composed of close-packed, hexagonal crystals, said crystals having a preferred orientation in at least the surface portions of said metal sheet with their basal planes in said surface portions, said metal sheet having a bi-axial ultimate strength at least 25% in excess of its uni-directional ultimate strength, said pressure vessel in its normal use subjecting said metal sheet to bi-axial tension stresses.

3. A pressure vessel composed at least in part of a rolled metal sheet, the metal of which comprises a titanium-base alloy containing about 2 to 10% aluminum, and characterized in having at atmospheric temperatures a predominantly alpha crystal structure composed of closepacked, hexagonal crystals, said crystals having a preferred orientation in at least the surface portions of said metal sheet with their basal planes in said surface portions, said metal sheet having a bi-axial ultimate strength at least 25% in excess of its uni-directional ultimate strength, said pressure vessel in its normal use subjecting said metal sheet to bi-axial tension stresses.

4. A pressure vessel composed at least in part of rolled sheet metal, said metal having a composition of about 7% aluminum, about 12% zirconium, balance titanium and exhibiting at atmospheric temperatures a crystal structure containing at least 90% by volume of a close-packed hexagonal alpha phase, the crystals of which at least in the surface portions of said metal sheet have a preferred orientation with their basal planes in said surface portions of said sheet and parallel thereto, said sheet metal being characterized by a bi-axial ultimate strength about in excess of its uni-directional ultimate strength, said pressure vessel in its normal use subjecting said metal sheet to bi-axial tension stresses.

5. A pressure vessel composed at least in part of rolled sheet metal, said metal comprising a titaniumba-se alloy containing 2 to 8% aluminum and exhibiting at atmospheric temperatures a crystal structure containing at least by volume of a close-packed hexagonal alpha phase, the crystals of which at least in the surface p tiOIlS Of said metal sheet have a preferred orientation 13 with their basal planes in said surface portions of said sheet and parallel thereto, said sheet metal being characterized by a bi-axial ultimate strength at least 25% in excess of its uni-directional ultimate strength, said pressure vessel in its normal use subjecting said metal 5 sheet to bi-aXial tension stresses.

References Cited by the Examiner UNITED STATES PATENTS 2,777,768 1/1957 Busch et all 75-1755 10 14 OTHER REFERENCES UNITED STATES PATENT OFFICE CERTIFICATE OF CORRECTION Patent No 5,266,951 August 16, 1966 Milton Bernard Vordahl It is hereby certified that error appears in the above numbered patent requiring correction and that the said Letters Patent should read as corrected below.

Column 1, line 42, for "fluid," read fluid column 2, line 10, for "means" read mean H line 61, for "extension" read extensive column 5, line 14, for "road" read load columns 9 and 10, TABLE III, in the heading to the fourth column for "Tens." read Tens. Elong. column 9, line 66, for "Ti-SAl-lZZr" read Ti-7Al-l2Zr column 10, line 60, for "pre-buldged" read pre-bulged column 11, line 50, for "0.05" read 0,5

Signed and sealed this 1st day of August 196 (SEAL) Attest:

EDWARD M.PLETCHER,JR. EDWARD J; BRENNER Attesting Officer Commissioner of Patents 

1. A PRESSURE VESSEL COMPOSED AT LEAST IN PART OF ROLLED METAL SHEET, THE METAL OF WHICH IS SELECTED FROM THE GROUP CONSISTING OF TITANIUM AND TITANIUM-BASE ALLOYS HAVING AT ATMOSPHERIC TEMPERATURES PREDOMONANTLY A CLOSE-PACKED, HEXAGONAL CRYSTAL STRUCTURE, THE CRYSTALS OF WHICH AT LEAST IN THE SURFACE PORTIONS OF SAID METAL SHEET HAVE A PREFERRED ORIENTATION WITH THEIR BASAL PLANES IN SAID SURFACE PORTIONS OF SAID SHEET AND PARALLEL THERETO, SAID METAL SHEET BEING CHARACTERIZED BY A BI-AXIAL ULTIMATE STRENGTH AT LEAST 25% IN EXCESS OF ITS UNIDIRECTIONAL ULTIMATE STRENGTH, SAID PRESSURE VESSEL IN ITS NORMAL USE SUBJECTING SAID METAL SHEET OT BI-AXIAL TENSION STRESSES. 